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摘要:
锂离子电池商用负极材料石墨比容量低，难以满足市场需求，金属有机骨架材料（metal-organic framework materials，MOFs）具有可调控的结构、较大的表面积和可调节的孔径，可用作下一代电化学储能器件，引起广泛研究。本文综述了金属（Fe、Co、Zn、Mn、Cu）基金属有机骨架及其衍生物的合成，重点介绍了以金属有机骨架材料为前驱体制备过渡金属氧化物（transition metal oxide，TMO）/C作为锂离子电池负极材料的研究进展，并对其发展方向进行了展望。
Abstract:The graphite as the commercial anode material for lithium-ion batteries shows the low specific capacity, which is difficult to meet the market demand. The metal-organic framework materials (MOFs) have the tunable structure, large surface area, and adjustable pore size, which can be used as the next generation of electrochemical energy storage devices, causing the extensive research. The synthesis of the metal (Fe, Co, Zn, Mn, Cu)-based metal organic frameworks and the derivatives were introduced in this paper, the research progress on the preparation of transition metal oxide (TMO)/C as the anode materials for lithium-ion batteries was focused, using MOFs as the precursors, and the development direction was prospected.
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稀有金属钼（Mo）是重要的高熔点金属，其熔点为2610 ℃，仅次于碳、钨、铼、钽和锇。金属Mo呈银白色，外形近似钢铁，具有高的硬度和弹性模量，低的蒸气压和蒸发速度，低的线膨胀系数，高的抗腐蚀能力等一系列优异特性，在现代国防、原子能工业、电真空、电光源等工程应用领域占有重要地位，在一些特殊高温应用领域甚至具有不可取代的作用[1–4]。
研究表明，在金属Mo基体中引入稀土氧化物粒子（可称作“稀土氧化物–Mo基材料”）可进一步提高材料的性能，拓展材料的应用。例如，通过引入氧化镧（La2O3）、氧化钇（Y2O3）等粒子对材料弥散强化，不仅可以大大提高金属Mo的室温强度和硬度，而且可以提高材料的再结晶温度，增强高温力学性能，显著延长作为高温发热体材料的使用寿命[5–8]。此外，在金属Mo基体中引入氧化钪（Sc2O3）、Y2O3等稀土氧化物粒子还可以提高材料的电子发射能力，用作优秀的阴极材料[9]。
作为改善金属Mo性能的稀土氧化物粒子，其尺寸大小及在Mo基体中的分布直接影响所制材料的性能。通常认为，粒子越细小，在Mo基体中分布越均匀，越有利于材料性能的提高[4–12]，因此，设法获得粒度细小的稀土氧化物粒子、并使其均匀分布在基体中，是制备高性能稀土氧化物–Mo基材料的基础。由于熔点较高，目前难熔金属主要采用粉末冶金方法制备，而在粉末冶金工艺中，原料粉末是决定材料性能和制造成本的关键一环，要获得高性能的稀土氧化物–Mo基材料，需要首先制备出高纯度、细粒度、稀土氧化物粒子细小且掺杂分布均匀的Mo基粉末原料。与传统制备稀土氧化物–Mo基粉末的机械合金化法相比，溶液燃烧法具有掺杂少、合成效率高、能耗低等优点。特别是溶液燃烧法的合成原料均为水溶性物质，目标金属在水溶液中以离子形态存在，能够很容易实现各组分在原子或分子水平上的均匀分散和混合，这为最终得到Mo基材料中稀土氧化物弥散相的粒径细化和均匀分布提供了有利条件。
为了增加溶液燃烧合成法的应用范围，同时为La2O3掺杂Mo合金的制备提供新思路，本文以七钼氨酸（(NH4)6Mo7O24·4H2O）作为金属源，甘氨酸（C2H5O2N）为燃料，硝酸铵（NH4NO3）为氧化剂，采用溶液燃烧法合成不同质量分数La2O3掺杂的Mo前驱体粉末，并对前驱体粉末进行还原、烧结，研究La2O3掺杂量（质量分数）对粉体性能及对烧结后Mo合金各项性能的影响。
1. 实验材料及方法
以高可溶性的七钼氨酸（(NH4)6Mo7O24·4H2O）为金属源，硝酸铵（NH4NO3）（≥99.0%）为氧化剂，甘氨酸（C2H5O2N）为燃料及添加剂，添加不同质量分数La(NO3)3·6H2O（以La2O3含量占最终合金粉末质量的比例为计算标准，分别为0、0.3%、0.7%、1.0%），通过溶液燃烧反应合成前驱体。在700 ℃下氢气氛围中还原，制备出La2O3掺杂Mo粉。对制备的粉末进行放电等离子体烧结（spark plasma sintering，SPS），烧结温度1600 ℃。
采用X射线衍射仪（X-ray diffraction，XRD；PANalytical X-Pert PRO MPD）对未添加La2O3的氧化钼前驱体及Mo–La2O3前驱体的物相组成进行表征。采用场发射扫描电子显微镜（field emission scanning electron microscope，FESEM；Hitachi SU8020）和透射电子显微镜（transmission electron microscope，TEM）对产物的显微组织进行观察。采用能谱仪（energy disperse spectroscope，EDS）对试样中Mo和La的元素分布进行测定。
2. 结果与讨论
2.1 La2O3掺杂量（质量分数）的Mo合金粉组织结构及性能的影响
图1为不同La2O3掺杂量的前驱体粉末微观形貌，可以清楚地发现，当不掺杂La2O3时，获得的前躯体粉末为片状结构，厚度为200 nm，片的尺寸约为0.5～2.0 μm。随着La2O3掺杂量的增加，其形貌开始变为细长颗粒状，且颗粒尺寸逐渐变小。当La2O3掺杂含量达到1.0%（质量分数）时，粉末晶粒尺寸以小于200 nm为主，且出现严重团聚现象。
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图 1 La2O3掺杂量对前驱体粉末显微形貌的影响：（a）0；（b）0.3%；（c）0.7%；（d）1.0%Figure 1. Effect of La2O3 doping content (mass fraction) on the microstructure of the precursor powders: (a) 0; (b) 0.3%; (c) 0.7%; (d) 1.0%对不同La2O3掺杂量的前驱体粉末在700 ℃下进行还原，图2为还原产物扫描电子显微形貌。由图可以看出，制备出的La2O3掺杂Mo粉尺寸在纳米级别，随着La2O3添加量的增加，Mo粉的晶粒尺寸逐渐减小，其中掺杂质量分数为0、0.3%、0.7%和1.0%La2O3的Mo粉晶粒尺寸分别为220、180、150以及100 nm，这是由于添加La2O3抑制了Mo晶粒长大。另外，由于纳米粉末尤其是难熔金属的纳米粉末的表面积非常大，为了降低体系能量，还原后的粉末颗粒自发的聚集在一起，从而出现了不均匀的团聚现象。
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图 2 掺杂不同质量分数La2O3的Mo粉700 ℃还原产物显微形貌：（a）0；（b）0.3%；（c）0.7%；（d）1.0%Figure 2. SEM images of the reduction products of the Mo powders doped by La2O3 in different mass fraction: (a) 0; (b) 0.3%; (c) 0.7%; (d) 1.0%图3为掺杂不同质量分数La2O3的Mo粉在700 ℃还原产物的X射线衍射图谱，由图可知，氧化钼前驱体均被还原成了Mo粉，这说明通过溶液燃烧法可以获得高纯度的La2O3掺杂Mo粉。此外，虽然在Mo粉中掺杂了不同含量的La2O3第二相粒子，但是在图中并未发现La的峰，可能是加入的La2O3所占比例非常小，在X射线衍射检测中未能发现。为了验证La2O3粒子的掺杂，实验对还原后的粉末进行了能谱分析，结果如图4所示，在掺杂质量分数为1.0%La2O3的Mo粉中发现了La特征峰，证明了La元素的存在。
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图 3 掺杂不同质量分数La2O3的Mo粉700 ℃还原产物X射线衍射图谱Figure 3. XRD patterns of the Mo powders doped by La2O3 in different mass fraction after reduction at 700 ℃![]()
图 4 掺杂质量分数1.0%La2O3的Mo粉在700 ℃还原产物的扫描电子显微形貌（a）和对应的能谱分析（b）Figure 4. SEM image (a) and the corresponding EDS analysis (b) of the Mo powders doped by 1.0%La2O3 after reduction at 700 ℃对还原后的粉末做进一步分析，通过透射电子显微镜对掺杂质量分数0.7%La2O3的Mo粉进行表征，结果见图5。从图中可以清楚地观察到，还原后的粉末粒径大约为150～200 nm，而且分散性较好。这主要是因为溶液燃烧法在反应过程中产生的前驱体晶粒细小，团聚体中存在大量的孔隙（如图1所示），因此在较低温度还原后，合金粉末的晶粒能够保持在纳米尺寸且分散性较好[13]。
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图 5 Mo–0.7La2O3前驱体粉末透射电子显微镜照片：（a）低倍；（b）高倍Figure 5. TEM images of the Mo–0.7La2O3 precursor powders: (a) low magnification; (b) high magnification2.2 La2O3掺杂量（质量分数）对Mo–La2O3烧结体组织结构及性能的影响
图6为经1600 ℃烧结后La2O3掺杂Mo合金的断口形貌。和纯Mo相比，La2O3掺杂Mo合金材料的晶粒更为细小，并且随La2O3质量分数的提高，细化作用逐渐明显。可以看出，在La2O3质量分数为0.7%时，Mo晶粒尺寸为500 nm左右，继续增加La2O3质量分数至1.0%，其晶粒尺寸降至300 nm。随着La2O3掺杂量的增加，Mo–La2O3烧结体中空隙数量增加，La2O3质量分数为1.0%时，其断口形貌中孔隙数量最多。
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图 6 经1600 ℃烧结后不同质量分数La2O3掺杂Mo合金的断口形貌：（a）0；（b）0.3%；（c）0.7%；（d）1.0%Figure 6. Fracture morphology of the Mo alloys doped by La2O3 in different mass fraction sintered at 1600 ℃: (a) 0; (b) 0.3%; (c) 0.7%; (d) 1.0%图7所示为不同La2O3掺杂量对Mo–La2O3合金相对密度的影响。可以明显看出，随着La2O3质量分数的提高，Mo合金的相对密度逐渐减小。这一方面是因为La2O3的实际密度低于纯Mo，随着掺杂量的提高，其相对密度必然会下降；另一方面，La2O3的加入会阻碍晶粒与烧结颈长大，同时阻碍晶界的迁移，使得材料的致密化行为变得困难，降低其相对密度[14]。这也与图6（d）中大量空隙相对应。
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图 7 1600 ℃烧结Mo–La2O3合金相对密度随La2O3质量分数变化Figure 7. Relative density of the Mo–La2O3 alloys doped by La2O3 in different mass fraction sintered at 1600 ℃图8所示为Mo–La2O3合金材料的显微硬度随着La2O3掺杂量的变化。从图中可以看出，合金材料的显微硬度呈现先增加后减小的趋势，在La2O3质量分数为0.7%时，显微硬度达到最高，为HV0.2546。这是由于La2O3的加入会阻碍晶粒生长，细化晶粒，提高材料的力学性能[15]。同时，第二相粒子La2O3可以起到钉扎作用，阻碍位错的迁移，使得材料硬度提高。但是，当La2O3掺杂量过多时，样品密度降低，孔隙数量增加，从而引起硬度降低[15‒16]。因此当La2O3掺杂量超过0.7%时，硬度值又出现下降的趋势。
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图 8 1600 ℃烧结Mo–La2O3合金显微硬度随La2O3质量分数变化Figure 8. Microhardness of the Mo–La2O3 alloys doped by La2O3 in different mass fraction sintered at 1600 ℃3. 结论
（1）将溶液燃烧法应用于纳米稀土氧化物掺杂Mo基材料的制备，成功制备出La2O3掺杂Mo合金粉，并经烧结获得合金样品，所制备合金样品具有优异的力学性能。
（2）随着La2O3掺杂量（质量分数）的增加，溶液燃烧合成制备的前驱体粉末逐渐由片状大颗粒变成细小的不规则颗粒。在掺杂量为1.0%时，前驱体粉末晶粒尺寸在200 nm左右。经还原后得到的Mo–La2O3粉末晶粒尺寸随着La2O3掺杂量的增加而减小，在掺杂量为1.0%时，晶粒尺寸为100 nm左右。
（3）所制得的La2O3掺杂Mo粉经1600 ℃烧结后产物相对密度在均在95%以上，随着La2O3掺杂量的增加（La2O3质量分数在0～1.0%范围内），相对密度逐渐降低，而显微硬度呈现先上升后下降的趋势。在La2O3掺杂量为0.7%时，Mo–La2O3合金显微硬度呈现出最大值，此时晶粒尺寸为500 nm左右，显微硬度达到HV0.2564。
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图 1 前驱体GO/ZIF-67（a）和G/Co3O4复合材料（b）显微形貌、G/Co3O4和Co3O4电极在100 mA·g−1电流密度下前三圈放电/充电曲线（c）以及在200 mA·g−1的循环图（d）[17]
Figure 1. Microstructures of GO/ZIF-67 precursor (a) and G/Co3O4 (b), the discharge/charge curves of G/Co3O4 and Co3O4 at the current density of 100 mA·g−1 for the first three cycles (c), and the cycling properties of G/Co3O4 and Co3O4 at 200 mA·g−1 (d)[17]
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图 3 空心多孔ZnO/C制备工艺流程（a）、空心多孔ZnO/C显微形貌（b）以及在电流密度为100 mA·g−1时空心多孔ZnO/C、空心多孔ZnO和商用ZnO的循环性能（c）[20]
Figure 3. Schematic diagram of hollow porous ZnO/C preparation process (a), microstructure of hollow porous ZnO/C (b), and cycling properties of hollow porous ZnO/C, hollow porous ZnO, and commercial ZnO at 100 mA·g−1 (c)[20]
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表 1 MOFs衍生锂离子电池负极材料
Table 1 MOFs-derived anode materials for lithium-ion batteries
MOFs前驱体 产物 电流密度 / (mA·g−1) 可逆容量 / (mA·h·g−1) 循环次数 参考文献 Co-MOF Co3O4/C 200 1052 60 [15] Co-MOF Co3O4/C 1000 601.0 500 [16] ZIF-67 G/Co3O4 200 714.0 200 [17] Fe-MOF C-Fe3O4 100 975.0 50 [18] Fe-ZIF Fe2O3@N-C 100 861.0 100 [19] MOF-5 ZnO/C 100 750.0 100 [20] Ppy-ZIF-8 ZnO/C 250 526.0 500 [21] Zn-BTC ZnO/C 100 919.0 100 [22] ZnO@ZIF-8 ZnO/C 2000 351.0 — [23] MOF-5 ZnO@C/CNT 100 758.0 100 [24] Mn-BTC MnO@C 3825 596.3 1000 [26] Mn-PBI MnO/C−N 300 1085.0 100 [27] Mn-BDC MnO/C@rGO 100 1536.4 100 [29] Cu-MOF CuOx-rGO 200 1490.0 220 [31] Cu-BTC CuO@C 100 1024.0 100 [32] [Cu(BTC)2]n-MOF CuO/C 100 510.5 200 [33] Cu-MOF CuO/C 100 789.0 200 [34] Cu-MOF CuO@C 1000 410.0 1000 [35] MIL-125@ZIF-67 Co3O4/TiO2 1000 838.6 600 [36] ZIF-67 Co3O4@Co3V2O8 100 948.0 100 [37] ZIF-67 Co3O4@TiO2 100 1057.0 100 [38] MOF-74-FeCo Co3O4-CoFe2O4 100 940.0 80 [39] ZnCo-MOF ZnO/ZnCo2O4/C 500 669.0 250 [40] GO/Zn-Co-ZIF/Ni rGO/ZnCo2O4-ZnO-C/Ni 100 1184.4 150 [42] 
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